Letter pubs.acs.org/JPCL
Electric Potential Gradient at the Buried Interface between LithiumIon Battery Electrodes and the SEI Observed Using Photoelectron Spectroscopy Julia Maibach,*,† Fredrik Lindgren,† Henrik Eriksson,† Kristina Edström,† and Maria Hahlin‡ †
Department of Chemistry − Ångström Laboratory, Uppsala University, Box 538, SE 751 21 Uppsala, Sweden Department of Physics and Astronomy, Uppsala University, Box 516, SE 751 21 Uppsala, Sweden
‡
S Supporting Information *
ABSTRACT: The buried interface between the bulk electrode material and the solid electrolyte interphase (SEI) in cycled Li-ion battery anodes is suggested to incorporate an electric potential gradient. This suggestion is based on photoelectron spectroscopy (PES) results from different anode materials that all show relative binding energy shifts between the components of the SEI and the active anode. Implications of this electric potential gradient on binding energy reference points in PES as well as on charge-transfer kinetics in Li-ion batteries are discussed. Specifically, we show that the separation of surface layer and bulk material spectral contributions (depth profiling) is crucial for consistent data interpretation. We conclude that previous interpretations of lithiation as cause for changes in PES spectra may need to be revised.
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bulk underneath the interface relative to those at the surface above the interface. In our ongoing work, synchrotron-based soft- and hard-X-ray photoelectron spectroscopy are frequently applied to investigate, for example, SEI and SPI layer formation and lithiation/ delithiation effects on the electrode materials as a function of layer thickness.14,16 The possibility to individually probe surface (i.e., the SEI/SPI) contributions using low photon energies as well as simultaneously probing surface and bulk (i.e., from the electrode material) specific spectral contributions using hard Xray PES (HAXPES) is one advantage of these synchrotronbased techniques. It is thus on the one hand possible to clearly identify SEI contributions and on the other hand simultaneously gain information about the composition and interaction of the SEI and the bulk active material; however, through recent HAXPES results we have come to realize that the character of the buried interface between the SEI and the bulk can strongly influence the assignments of PES peaks. If these interfacial effects are not accounted for, it will alter the data interpretation, leading to faulty SEI compositions and overinterpretation of chemical shifts caused by lithiation. The initial SEI formation on clean and well-defined surfaces has, for example, been demonstrated in UHV studies for lithium cobalt oxide (LCO)17 and graphite18 electrode materials. In these results electric potential gradients are observed between surface layers and bulk materials; however,
he interfaces in lithium-ion batteries (LIBs) govern essential device properties such as safety, lifetime, and charging kinetics. Therefore, characterizing and understanding the interplay between electrolyte and electrode materials in LIBs is crucial for the development of improved battery systems. Upon contact and during the first cycle new phases are formed at the interface between the electrodes and the electrolyte. Negative electrodes form the so-called solid electrolyte interphase (SEI)1−3 and positive electrodes form the “solid permeable interphase” (SPI).4 The ideal SEI passivates the anode surface to avoid continuous electrolyte decomposition, which makes it a key factor for battery stability and safety. The SEI is generally composed of inorganic and organic, polar, and nonpolar molecules that solidify at the interface during cycling, with thicknesses in the range of a few tens of nanometers. Likewise but thinner (3 to 4 nm), the SPI is formed by electrolyte decomposition as well as dissolution products of active materials and affects the aging of a battery.5 In-depth studies of SEI formation are an ongoing topic in LIB research; the SEI has been investigated for many material combinations and with a large variety of techniques,3,6−12 and photoelectron spectroscopy (PES) has been proven to be one of the most effective.13,14 This is because PES probes surfaces and interfaces with high surface sensitivity, and it is one of the few techniques that can probe such thin layers.15 We consider the influence of the electrode state of charge, the SEI itself, as well as the interface between SEI and electrode on relative photoelectron binding energies. Consequently, we discuss consistent assignment of chemical species present in the © 2016 American Chemical Society
Received: February 19, 2016 Accepted: April 22, 2016 Published: April 22, 2016 1775
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top of Figure 1, peaks from the carbon coating and carbon black additive (CC/CB, gray) and the binder PVdF (CH2 at 286.4 eV, cyan and CF2 at 290.5 eV, blue) are observed. Also shown are the spectra for the first insertion of 0.5 Li per formula unit, complete lithiation and delithiation, as well as the second complete lithiation. The corresponding electrochemical cycling curves are shown in the SI in Figure S1. For the cycled electrodes, the carbon black peak (gray), which is attributed to the bulk electrode, is still clearly visible. Additionally, several new components due to SEI formation (various shades of blue) such as hydrocarbon, ether compounds, and carbonates appear. All spectra in Figure 1a,b are referenced in binding energy to the surface components; the pristine sample and the sample after insertion of 0.5 Li per formula unit are calibrated towards the CF2 binder component at 290.5 eV, and the remaining spectra are calibrated to the hydrocarbon peak at 285 eV. One can clearly observe a significant but largely reversible change in binding energy position for the bulk carbon black peak (and the Ni0 peak in Figure 1 b) depending on the state of lithiation, while the components attributed to the surface (i.e., SEI) align perfectly. This same alignment is also true for the other elemental lines (F 1s, O 1s, P 2p, and Li 1s) originating from the SEI (see ref 20), while bulk components (Ni 2p, Ti 2p) vary in position. When plotting the same data but using the carbon black emission for binding energy referencing, all bulk-electrode material peaks align as is exemplified in Figure 1c,d for C 1s and Ni 2p, while the surface components (SEI peaks) vary in binding energy depending on the state of electrode lithiation. From this we draw the conclusion that a change in relative binding energy between bulk and surface occurs, which differs in magnitude depending on the potential at which the cycling was stopped. Clearly, the change becomes most pronounced between potentials where SEI formation takes place (see Figure S1). This is evident in the Ni0.5TiOPO4 example (Figure 1 a) when comparing the 0.5 Li/formula unit spectra to the fully lithiated spectra. In the former, only a minor variation in the CC/CB binding energy is observed, while for the fully lithiated material after SEI formation a very pronounced shift of nearly 1 eV is observed. The connection to the onset of SEI formation becomes even clearer when looking at spectra from graphite electrodes cycled in half-cells versus Li metal shown in Figure 2. At OCV, the surface hydrocarbon and the bulk graphite peaks are separated by 0.7 eV, which is in accordance to literature.21 At the beginning of SEI formation (i.e., ∼0.8 V vs Li+/Li), a shift in the bulk graphite peak relative to the growing surface hydrocarbon peak can be observed. This relative shift increases to 1.8 eV when lowering the potential to 0.4 V versus Li+/Li, which is commonly considered higher than the graphite lithiation potential (starting at 0.2 V vs Li+/Li22,23). At the end of lithiation at 0.01 V the overall shift of bulk graphite versus the surface hydrocarbon is 2.65 eV, out of which we attribute only a minor fraction to a chemical shift caused by the lithiation of graphite. This is of particular importance in the case of Ni0.5TiOPO4, where lithiation should not be considered as the major cause for the binding energy variations between bulk and surface as the lithiation of the electrode material can be followed by the change of oxidation state from Ni2+ to Ni0 in the Ni 2p component (see Figure 1d); however, no further shift due to lithium-ion (de)insertion is observed as the Ni0 component stays constant in binding energy (relative to the CC/CB peak) for all cycled electrodes.
the influence of this potential gradient in PES on cycled electrodes as well as on electrochemical processes and reaction kinetics is so far not widely considered in the battery community but may have significant implications on chargetransfer kinetics in a working battery. We present a model based on an interfacial potential gradient to explain experimental data and raise awareness of the importance of correct binding energy referencing for cycled battery electrodes, as this is a current topic in PES on battery materials.19 We exemplify the buried interface’s influence with PES data from previous work on carbon-coated nickel titanium oxyphosphate (Ni0.5TiOPO4) (ref 20) and graphite. The relatively new anode material Ni0.5TiOPO4 is a particularly interesting system to elucidate the buried interface’s effects because it forms an SEI consisting of poorly conductive species, and at the same time, the lithiation process can be followed in core-level spectra (such as Ni or Ti) that are not involved in the SEI formation. Additionally, the carbon black additive present in the bulk composite electrode should not take part in the lithiation reaction at the applied potentials. Thus, the effects of lithiation on photoelectron binding energies can be well separated from interfacial effects. In Figure 1a,b, the C 1s and Ni 2p HAXPES spectra (incoming photon energy = 2300 eV) of a Ni0.5TiOPO4 composite electrode (75 wt % Ni0.5TiOPO4/15 wt % Super P carbon/10 wt % polyvinylidine fluoride (PVdF)) are shown, respectively. In the spectra of the pristine sample shown in the
Figure 1. (a) C 1s spectra and (b) Ni 2p spectra of Ni0.5TiOPO4 electrodes at different state of lithiation/delithiation referenced in binding energy to surface hydrocarbon. (c) C 1s and (d) Ni 2p spectra of Ni0.5TiOPO4 electrodes at different states of lithiation/delithiation referenced in binding energy to the bulk carbon black signal. 1776
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Figure 2. Bulk (black) and surface-sensitive (red) C 1s spectra of graphite electrodes at different potentials during the first lithiation referenced in binding energy to the hydrocarbon peak at 285 eV.
Our data show that lithiation can be excluded as the major reason for the relative binding energy shifts between SEI and bulk electrode components. Differential charging or an extended potential distribution within the SEI layer also seem unlikely to cause the relative binding energy variations as both phenomena should lead to asymmetric peak shapes,24−26 which is not the case for the presented data. Further evidence against differential charging was obtained by comparing spectra as a function of incoming photon intensity, and no changes in the relative peak positions or line shape as a function of intensity were observed (see SI, Figure S2). Therefore, alternative explanations have to be considered. We believe that the origin of this deviation in bulk binding energies with respect to surface components is located at the buried interface between the electrode material and the SEI. One possibility to understand the physical origin of such a change in binding energy between surface and bulk would be the presence of an electric potential gradient at the interface (e.g., in form of a dipole layer, see Figure 3), as has been postulated previously in literature both in experimental18,19,27 and theoretical28 work within the battery context. In PES, this potential gradient is generally quantified directly from work function changes across the interface of the two phases in contact. The difficulty with ex situ cycled battery electrodes is that the work function at least for the buried bulk component is not directly accessible; however, based on what we know from electrochemistry, bulk electrode and SEI have different electrochemical potentials (which may be interpreted as the Fermi level according to ref 29) and the SEI has ideally no electronic conductivity as the SEI passivates the surface, thereby hindering continuous electrolyte reduction. If this is the case, then bulk electrode and surface layer would not be on the same Fermi level with respect to the PES analyzer similar to the observed Fermi level decoupling for poorly conducting layers on conductive substrates.26,30 However, the binding energy in PES is generally referenced to the Fermi level (of the analyzer). For conductive samples this is straightforward as the Fermi levels in the sample and the analyzer are well-defined and can be measured with PES.15 For poorly or nonconductive samples (such as battery electrodes with an SEI), a different calibration approach is usually applied based on an internal reference. From polymer films the usage of
Figure 3. Model for relative binding energy differences based on changed interface dipole strength depending on state of lithiation.
adventitious or any other source of saturated hydrocarbon as an internal point of binding energy reference31,32 was adapted to batteries. These hydrocarbons are typically located at the sample surface, and thus their binding energy position might not be representative for underlying (conductive) bulk materials. Both the Fermi level and surface calibration methods of referencing are valid in the respective case and also if the poorly conductive surface layer thickness is larger than the PES probing depth; however, if the probing depth exceeds the thickness of the surface layer and components of the wellconducting bulk material appear in the same spectra as surface species (which often is the case of HAXPES data on battery electrodes), the observed binding energy shift cannot be interpreted as a regular chemical shift. Thus, contributions to the observed relative binding energy shifts between SEI and bulk electrode components appear to originate to a large extent from different physical binding energy reference points for surface and bulk. In our model, these different binding energy reference points are linked to the presence of an electric potential gradient between surface layer and bulk. If one compares the three bottom spectra in Figure 1 (first fully lithiated, fully delithiated, second fully lithiated), one can clearly see that the relative binding energy shift between bulk and surface changes with the electrode’s state of charge but at this stage is fairly independent of SEI formation (as SEI formation mainly occurs during the first lithiation, see Figure S1). The potential gradient at the buried interface between the two components thus changes with the electrode’s state of charge. In the presented example, a large binding energy difference between bulk carbon black and the SEI hydrocarbon is 1777
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Furthermore, it is emphasized that a one-point binding energy calibration based on surface components might lead to misinterpretation of the data and faulty assignment of chemical species to certain peaks in PES Li-ion battery characterizations. On the basis of the observations described here, we believe that the previous attribution of shifts in the bulk-originating signals such as graphite or carbon black to lithiation effects37 might need to be revised.
observed for the fully lithiated Ni0.5TiOPO4. Applying our interfacial potential gradient model to this data set, we explain the binding energy difference (ΔEB) by a large double-layer potential drop caused by a high amount of accumulated charges at the buried interface (see Figure 3 a). For the delithiated state the binding energy difference between bulk and surface component is reduced; this is depicted as a reduced potential drop (i.e., reduced amount of interface charge or reduced number of dipoles) in Figure 3 b; however, this reduced shift could also be explained by a changed orientation of the electric field with respect to the surface normal. The presented model, based on interfacial effects, can also be applied to other battery material systems such as semiconducting silicon electrodes33 and Na-ion batteries34 in which a similar trend upon surface layer formation during cycling was observed. Considering battery operation, it should be noted that the presence of a double-layer potential drop at the interface between electrode and SEI could influence the (de)lithiation kinetics of various electrode materials (as anticipated, for example, in refs 28 and 35) because it may act as a barrier for lithium-ion insertion/extraction depending on the dipole orientation; however, in the full device, also the interaction of the SEI surface with the electrolyte needs to be considered because a mirror potential gradient at the solid−liquid interface could form. Thus, in situ experiments in which the contact between an electrode material and various SEI components is built up through stepwise evaporation and successive PES characterization as well as ambient pressure PES measurements of the solid−liquid interface would be highly beneficial. From a chemical perspective the interface potential gradient could be caused by polar components such as adsorbed/ reduced solvent molecules,18,28 LiF, and so on, which are formed during SEI build-up and, for example, commonly used to modify electronic energy level alignment in organic electronics.36 From an electrochemical perspective, also a two-phase capacitive interface storage mechanism as described by Fu et al.35 can lead to the observed potential drop. This would also be consistent with a previous report on a high concentration of Li-ions leading to nonstoichiometric LixC compounds at the buried interface between the SEI and lithiated graphite.37 Possible physical factors contributing to the potential gradient could be the surface layer thickness and partial dissolution of polar or charge-carrying SEI components during cycling (corresponding to previously observed SEIthickness variations upon cycling7,8). Additionally, the potential at which the respective cycling was stopped (equivalent to a certain degree of lithiation) before retrieving the electrode for PES analysis plays a role, as this influences the electrode’s Fermi level; however, trying to separate these factors will be the subject of a future study. We present PES data of cycled battery electrodes where relative binding energy shifts are observed when surface layer formation takes place. We discuss a model based on interfacial potential gradients leading to different physical points of binding energy reference in PES to explain these shifts. The presented examples show that it is crucial to perform both highly surface-sensitive as well as more bulk-sensitive PES measurements on cycled battery electrodes to obtain a clear picture of which species are located in the surface film and which originate from the bulk electrode because binding energy differences due to interfacial effects can occur.
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EXPERIMENTAL METHODS All experiments were performed on composite electrodes, consisting of combinations of highly conductive components with metallic character such as the graphite and carbon black additive, more semiconductor-type components such as the oxyphosphate, and even insulating part such as the binder. The experimental details of electrode fabrication and cycling for the Ni0.5TiOPO4 study have been published in detail in ref 20. The photoelectron spectroscopy measurements of these samples were performed at the KMC-1 beamline using the HIKE end station at BESSY II (Helmholtz Zentrum, Berlin, Germany). The spectra were recorded using first-order light at 2300 eV monochromatized using the Si(111) crystal pairs. The Scienta R4000 electron analyzer was set to a pass-energy of 500 eV. The graphite electrodes were prepared from a slurry containing 85 wt % Graphite (Hitachi), 3 wt % Graphite KS6 (Timcal), 2 wt % carbon black powder, and 10 wt % Kynar binder (Arkema) in N-methyl-2-pyrrolidone. The slurry was mixed with a planetary ball mill for 1 h and then coated on a copper foil (dry coating thickness was 30 μm.) Coffee-bag battery (aluminum pouch) cells were prepared with graphite as working electrode, Solupor separator, and Li-foil as counter electrode. 1 M LiPF6 in a solvent mixture of ethylene carbonate (EC) and diethyl carbonate (DEC) with a volume ratio 2:1 was used as electrolyte. The graphite electrodes were lithiated to different potentials in the first cycle with a CCCP (constant current constant potential) program on a battery tester from Arbin. Constant current was calculated to correspond to a cycling rate of C/10. All batteries were allowed to stabilize after cycling during a constant voltage step until the current was lower than the current for C/100. XPS sample preparation and transfer followed ref 38. All samples were mounted on sample holders using conductive Cu-tape. PES characterization of the graphite electrodes was performed at the soft X-ray beamline I411 at the MAX IV laboratory, Lund, Sweden. PES spectra were obtained using photon energies such that the kinetic energy of the emitted photoelectrons was 140 (probing depth 2 nm) and 550 eV (probing depth 7 nm).14
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ASSOCIATED CONTENT
* Supporting Information S
The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acs.jpclett.6b00391. Electrochemical cycling data of Ni0.5TiOPO4 (Figure S1). Example for data collection as a function for different incoming photon intensities (Figure S2). (PDF)
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AUTHOR INFORMATION
Corresponding Author
*E-mail:
[email protected]. 1778
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Soft and Hard X-Ray Photoelectron Spectroscopy. Electrochim. Acta 2013, 97, 23−32. (15) Hüfner, S. Photoelectron Spectroscopy: Principles and Applications; Springer: New York, 2003. (16) Högström, K. C.; Hahlin, M.; Malmgren, S.; Gorgoi, M.; Rensmo, H.; Edström, K. Aging of Electrode/Electrolyte Interfaces in Lifepo4/Graphite Cells Cycled With and Without PMS Additive. J. Phys. Chem. C 2014, 118, 12649−12660. (17) Hausbrand, R.; Cherkashinin, G.; Ehrenberg, H.; Gröting, M.; Albe, K.; Hess, C.; Jaegermann, W. Fundamental Degradation Mechanisms of Layered Oxide Li-Ion Battery Cathode Materials: Methodology, Insights and Novel Approaches. Mater. Sci. Eng., B 2015, 192, 3−25. (18) Lee, C.; Mun, B.; Ross, P. N. The Chemical Reaction of Diethyl Carbonate with Lithium Intercalated Graphite Studied by X-Ray Photoelectron Spectroscopy. J. Electrochem. Soc. 2002, 149, A1286− A1292. (19) Oswald, S. Binding Energy Referencing for Xps in Alkali MetalBased Battery Materials Research (I): Basic Model Investigations. Appl. Surf. Sci. 2015, 351, 492−503. (20) Eriksson, R.; Lasri, K.; Gorgoi, M.; Gustafsson, T.; Edström, K.; Brandell, D.; Saadoune, I.; Hahlin, M. Electronic and Structural Changes in Ni0.5TiOPO4 Li-Ion Battery Cells Upon First Lithiation and Delithiation, Studied by High-Energy X-Ray Spectroscopies. J. Phys. Chem. C 2015, 119, 9692−9704. (21) Barr, T. L.; Seal, S. Nature of the Use of Adventitious Carbon as a Binding-Energy Standard. J. Vac. Sci. Technol., A 1995, 13, 1239− 1246. (22) Dahn, J. R.; Fong, R.; Spoon, M. J. Suppression of Staging in Lithium-Intercalated Carbon by Disorder in the Host. Phys. Rev. B: Condens. Matter Mater. Phys. 1990, 42, 6424−6432. (23) Ohzuku, T.; Iwakoshi, Y.; Sawai, K. Formation of LithiumGraphite Intercalation Compounds in Nonaqueous Electrolytes and Their Application as a Negative Electrode for a Lithium Ion (Shuttlecock) Cell. J. Electrochem. Soc. 1993, 140, 2490−2498. (24) Bryson, C. E. Surface-Potential Control in XPS. Surf. Sci. 1987, 189, 50−58. (25) Tielsch, B. J.; Fulghum, J. E. Differential Charging in XPS.1. Demonstration of Lateral Charging in a Bulk Insulator Using Imaging XPS. Surf. Interface Anal. 1996, 24, 422−427. (26) Yu, X. R.; Hantsche, H. Vertical Differential Charging in Monochromatized Small Spot X-Ray Photoelectron-Spectroscopy. Surf. Interface Anal. 1993, 20, 555−558. (27) Tonti, D.; Pettenkofer, C.; Jaegermann, W. Origin of the Electrochemical Potential in Intercalation Electrodes: Experimental Estimation of the Electronic and Ionic Contributions for Na Intercalated into TiS2. J. Phys. Chem. B 2004, 108, 16093−16099. (28) Leung, K.; Leenheer, A. How Voltage Drops Are Manifested by Lithium Ion Configurations at Interfaces and in Thin Films on Battery Electrodes. J. Phys. Chem. C 2015, 119, 10234−10246. (29) Reiss, H. The Fermi Level and the Redox Potential. J. Phys. Chem. 1985, 89 (18), 3783−3791. (30) Barr, T. L. Studies in Differential Charging. J. Vac. Sci. Technol., A 1989, 7, 1677−1683. (31) Beamson, G.; Briggs, D. High Resolution XPS of Organic Polymers: The Scienta Esca300 Database; John Wiley & Sons, Ltd.: West Sussex, England, 1992. (32) Siegbahn, K.; Nordling, C.; Fahlman, A.; Nordberg, R.; Hamrin, K.; Hedman, J.; Johansson, G.; Bergmark, T.; Karlson, S.-E.; Lindgren, I. ESCA; Atomic, Molecular and Solid State Structure Studied by Means of Electron Spectroscopy; Nova acta Regiae Societatis Scientiarum Upsaliensis 20; Almqvist & Wiksells: Uppsala, Sweden, 1967. (33) Lindgren, F.; Xu, C.; Maibach, J.; Andersson, A. M.; Marcinek, M.; Niedzicki, L.; Gustafsson, T.; Björefors, F.; Edström, K. A Hard XRay Photoelectron Spectroscopy Study on the Solid Electrolyte Interphase of a Lithium 4,5-Dicyano-2-(Trifluoromethyl)Imidazolide Based Electrolyte for Si-Electrodes. J. Power Sources 2016, 301, 105− 112.
The authors declare no competing financial interest.
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ACKNOWLEDGMENTS We are very thankful to Ismael Saadoune (Uni. Cadi Ayyad, Marrakech, Morocco) for the supply of the Ni0.5TiOPO4 material for the original study. We thank HZB for the allocation of synchrotron radiation beamtime. Additionally, we gratefully acknowledge experimental assistance from the staff at Max IV Laboratory, Lund, Sweden. The research leading to these results has received funding from the European Union’s Seventh Framework Programme (FP7/2007-2013) under grant agreement no. 608575. Further financial support from StandUp for Energy and the Swedish Research Council (grant VR-2012-4681) is gratefully acknowledged.
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